Materials Transactions, Vol. 43, No. 10 (2002) pp. 2400 to 2407Special Issue on Superplasticity and Its Applicationsc©2002 The Japan Institute of Light Metals

Continuous Dynamic Recrystallization in a Superplastic7075 Aluminum Alloy

Xuyue Yang, Hiromi Miura and Taku Sakai

Department of Mechanical Engineering and Intelligent Systems, The University of Electro-Communications,Tokyo 182-8585, Japan

New grain evolution taking place during superplasticity was studied by means of tensile tests as well as metallographic observationfor a unrecrystallized coarse-grained 7075 aluminum alloy. Grain boundary sliding (GBS) frequently takes place even on the layered highangle boundaries (HABs) parallel to the tensile axis and brings about rotation of subgrains near the HABs and subsequently in grain interiors.The misorientations of (sub)grain boundaries evolved in the pancaked grains increase accompanied by a randomization of the initial texture,followed by development of new grains with HABs. This indicates that unrecrystallized and pancaked grain structure developed by cold rollingis an important prerequisite not only for the appearance of superplasticity, but also for the dynamic evolution of new fine grains. It is concludedthat the mechanism of new grain evolution can be a deformation-induced continuous reaction, that is continuous dynamic recrystallization(CDRX). A model for CDRX is discussed in detail comparing with previous several models.

(Received April 15, 2002; Accepted July 30, 2002)

Keywords: 7075 aluminum alloy, superplasticity, pancaked grain, grain boundary sliding, grain rotation, misorientation, continuousdynamic recrystallization

1. Introduction

Some pseudo-single aluminum alloys exhibit high super-plasticity when they are processed into a fine-grained struc-ture below 10 µm either by static or dynamic recrystallizationthrough appropriate thermo-mechanical treatments.1, 2) Underdynamic condition, fine recrystallized grain structures can bedeveloped in as-cold or warm-worked materials by continu-ous reaction during early stages of hot deformation. Afterthe works of Watts et al.,3) the fine-grained structures dynam-ically evolved in unrecrystallized aluminum alloys has beenstudied and this process is often termed as continuous dy-namic recrystallization (CDRX).1–7) Several mechanisms op-erating during CDRX have been proposed for dynamic for-mation of high angle boundaries (HABs), e.g. long-range mi-gration of subboundary or subgrain growth,4, 5) pile-up andabsorption of dislocations into subboundaries6) and subgrainrotation7) etc. It is still unclear, however, how low angle sub-grain boundaries evolved in grain interiors transform to highangle ones under hot deformation.

The authors7–12) have studied superplasticity taking placein a cold-rolled 7075 aluminum alloy which was previ-ously warm-deformed just before hot deformation. The re-sults showed that the microstructures evolved under warm-deformation to over a critical strain can not be recrystallizedfor a long period of time (e.g. 104–106 s) at high tempera-tures,8) but leading to the appearance of fine grained struc-ture and superplasticity during hot deformation.9) The strainrate dependence of flow stresses at low strains is clearlychanged in the three regions of strain rate, although the lay-ered coarse-grained structure introduced by cold rolling ex-isted stably.9, 10) This suggests that grain boundary sliding cantake place even on the layered prior HABs parallel to the ten-sile axis.10–12) The structural mechanisms of fine grain evolu-tion during deformation, however, has been still unclear. Oneof the reasons may be in the pervious microstructural obser-

vation carried out mainly on the rolling plane.The aim of the present work is to study the evolution pro-

cesses of fine grained structure taking place in a unrecrystal-lized 7075 aluminum (Al) alloy and the character distributionof these new grain boundaries evolved during hot deforma-tion. The present work is especially concerned with the mi-crostructural evolution taking place in the longitudinal-shorttransverse section of a specimen. The microstructures thatare evolved under deformation were observed in-situ using ascanning electron microscope (SEM) equipped with a heatingand tensile stage. The SEM was incorporating an orientationimaging microscopy (OIM) technique obtained from electronback scattering diffraction (EBSD) patterns. The mechanismsof microstructural development and its relationship with su-perplasticity are discussed in detail.

2. Experimental Procedure

The material tested was a commercially produced 7075 Alalloy with the following chemical composition; Zn 5.62, Mg2.63, Cu 1.58, Cr 0.2, Fe 0.16, Si 0.07, Ti 0.02, Mn 0.05 andbalance Al (all in mass%). The as-received 3 mm thicknessplates of the 7075 Al alloy were machined to 2 to 1.33 mm andthen homogenized at 763 K for 10.8 ks in a molten salt bath,followed by quenching in water. The average grain size wasabout 17 µm. These plates were cold rolled to 1 mm thick-ness, leading to three kinds of reduction, i.e. 25%, 50% and67%. The specimens with a gauge section of 12 mm lengthand 5 mm width were machined parallel to the rolling direc-tion.

Tensile tests were conducted in vacuum at temperatureshigher than 773 K using an Instron-type testing machine,which was equipped with a hydrogen gas quenching appara-tus.8, 9) The specimens were deformed first to a strain of 40%at 623 K and at 4 × 10−3 s−1 because cold-rolled microstruc-tures were stabilized by η-phase precipitates.8) They were re-

Continuous Dynamic Recrystallization in a Superplastic 7075 Aluminum Alloy 2401

Fig. 1 A series of typical true stress vs. nominal strain curves at 773 K andat various strain rates for 7075 Al alloy with pancaked grains developed bycold rolling. The samples were pre-deformed to a strain of 40% at 623 Kand 4 × 10−3 s−1 just before hot deformation.

heated and kept for 1 ks at 773 K or 798 K and then deformedin tension at various initial strain rates from 2.9 × 10−5 s−1

to 2.1 × 10−1 s−1. They specimens were quenched by H2 gasimmediately after hot deformation, which quenched in hot de-formed microstructures and stopped precipitation taking placeduring cooling. The deformed microstructures were exam-ined mainly on the longitudinal section (L-ST) of deformedspecimens by using scanning electron microscopy (SEM).The measurement of new grain orientation and its bound-ary misorientation were performed by using EBSD technique.Some marker lines were scratched by using diamond paste onthe specimen surfaces parallel to the L-ST plane. The mor-phology of grain structures and grain boundary sliding as wellas the development of fine grains with their crystal orienta-tions were observed by SEM/OIM.

3. Experimental Results

3.1 Hot deformation behaviorA series of true stress-nominal strain (σ−ε) curves at 773 K

is shown in Fig. 1. Flow stresses as well as the shape of σ − ε

curves are sensitively dependent on strain rate. The σ − ε

curves exhibit a sharp stress peak just after yielding, followedinitially by a flow softening, a stress dip and then a flow hard-ening at high strains. The stress drop during flow softeningdecreases with decreasing strain rate, while a flow hardeningfollowed by a little softening appears at the lowest strain rate.

Strain rate dependence of the peak flow stress (σp) appear-ing just after yielding is represented in Fig. 2. It is evidentthat the strain rate dependence of σp changes clearly in thethree regions of strain rate, ε; i.e. the region I (lower ε), II(medium ε) and III (higher ε). The elongation to fracture arebeyond 300% at 773 K except the highest strain rate and wasover 750% at 793 K in the region II.9) It is considered, there-fore, that typical superplasticity takes place in the region IIaccompanied with a maximum total elongation and the stressexponent of about 1.7 in Fig. 2.

Figure 3 shows an effect of the reduction of cold rolling onthe σ −ε curves at 798 K and at 2.9×10−3 s−1. It can be seen

Fig. 2 Strain rate dependence of the peak flow stress (σp) at 773 K and798 K for 7075 Al alloy with pancaked grains.

Fig. 3 Effect of cold rolling reduction on true stress vs. nominal straincurve for 7075 Al alloy with pancaked grains. The samples werepre-deformed to a strain of 40% at 623 K and 4 × 10−3 s−1 just beforehot deformation.

here that the flow behavior is sensitively affected by rollingreduction, i.e. the total elongation increases and converselythe peak flow stress decreases with increase in the rolling re-duction. It was shown in the previous work12) that increasein the reduction results in decrease in the grain size in the STdirection (i.e. the thickness direction) (DST). Figure 3 sug-gests, therefore, that decrease in the grain size can accelerateappearance of superplastic flow behaviors.

3.2 Microstructural changes during deformationMicrostructural changes and new grain boundary evolution

taking place in the early stage of hot deformation were ob-served by SEM/OIM. Typical OIM micrographs for the sam-ples, just before deformation (ε = 0) and deformed to 160%at 798 K and at 2.9 × 10−3 s−1, are represented in Fig. 4.Different color levels indicate different crystallographic ori-entations. Here high-angle boundaries with misorientationsmore than 15◦ are delineated by thick-black lines, while low-angle boundaries in the range of 4◦ and 15◦ by thin-black linesand those in the range 2◦–4◦ by red lines. It can be seen inFigs. 4(a) and (c) that deformation bands with moderate an-gle boundaries as well as recovered subgrains are developed

2402 X. Yang, H. Miura and T. Sakai

Fig. 4 OIM micrographs of 7075 Al alloy with pancaked grains deformed at 798 K and at 2.9×10−3 s−1, followed by H2 gas quenching.Tensile axis was parallel to the pancaked layered grains (i.e. the rolling direction). Crystallographic orientations are expressed incolors according to inverse pole figure and color code is consistent with tensile axis. Thin-red lines correspond to boundaries ofmisorientation > 2◦, thin-black lines > 4◦ and thick-black lines > 15◦, respectively. (a), (c) ε = 0% and (b), (d) ε = 160%.

in coarse grain interiors in the L-LT plane (i.e. the rollingplane), while high density layered grains appear roughly par-allel to the tensile axis in the L-ST plane (i.e. the side surface).All of these high angle boundaries (HABs) is considered tobe the original ones existed just before hot deformation.10, 12)

Figures 4(a) and (c) also shows that low angle boundaries(LABs) below 4◦ as will as medium angle ones from 4◦ to 15◦are frequently developed in these pancaked grain interiors.The medium angle boundaries may be resulted from deforma-tion bands developed by prior-cold rolling.9) After straining to160%, such layered grain structures are fully destroyed andreplaced by an almost equiaxed, uniform fine-grained struc-ture with HABs, as can be seen in Figs. 4(b) and (d). Thenew grains evolved have a few LABs and also are randomlydistributed.

Figure 5 shows the inverse pole figures for the specimen de-formed to several strains at 798 K and at 2.9 × 10−3 s−1. Theinitial structure is oriented roughly on the line from 〈001〉 to〈111〉 which was developed by cold rolling.13) With increas-ing strain, this texture is progressively destroyed and finallyreplaced to a very diffused and nearly random one at a strainof 160%. These results in Figs. 4 and 5 suggest that the dy-namic evolution of new fine grains can result from (sub)grainrotation taking place frequently in the early stage of hot de-formation, leading to a rapid increase in misorientation.

The distribution of the misorientations above 2◦ for(sub)grain and dislocation boundaries developed in the wholearea is shown at strains of 0, 40 and 160% in Figs. 6(a), (b),and (c), respectively. For reference, the Mackenzie distribu-

tion14) for randomly oriented cubes are also shown in dashedline in Fig. 6. In the initial microstructure just before hot de-formation (i.e. ε = 0), the distribution shows a bimodal shape,where the peaks are at below 5◦ and at around 40◦. The formercorresponds to subgrain boundaries evolved during reheatingand the latter to the original layered HABs. The volume frac-tion of LABs decreases with increasing deformation, whilethat of HABs increases. At a strain of 160%, the misorienta-tion distribution shows a single peak type and approaches tothe Mackenzie distribution. The average value (θav) of 35.3◦is, however, rather smaller than that for the Mackenzie distri-bution, i.e. 41.2◦,14) because the frequency of low misorien-tations below 5◦ is still high. This should be a characteristicof the grain structure dynamically evolved, because LABs arealways introduced during hot deformation.

3.3 In-situ observation by SEM/OIMFigure 7 shows a series of SEM micrographs for the ex-

actly same place in a longitudinal section of the same sampledeformed to various strains at 798 K and at 2.9 × 10−3 s−1. Itcan be seen in Fig. 7(a) that layered grain boundaries appearclearly in the initial mirror-like surface even after deforma-tion to 10%. The grain boundary appearance itself suggeststhat grain boundary sliding (GBS) took place on the HABsof such layered grains, although they are roughly parallel tothe tensile axis. It is interesting to note that various portionsof such layered grain boundaries fluctuate and are frequentlyserrated, and then parts of them are not parallel to the tensiledirection (see also Fig. 10). With increasing strain, pancaked

Continuous Dynamic Recrystallization in a Superplastic 7075 Aluminum Alloy 2403

Fig. 5 Strain dependence of inverse pole figures for rolling direction of the plate of 7075 Al alloy with pancaked grains. The sampleswere deformed to several strains at 798 K and at 2.9 × 10−3 s−1. (a) ε = 0%, (b) ε = 40% and (c) ε = 160%.

0.0

0.1

0.2

0.3

0.0

0.1

0.2

0 10 20 30 40 50 60 700.0

0.1

0.2

Misorientation angle, / deg

Fre

quen

cy ,

f

7075Al = 0

= 160

Fre

quen

cy ,

f F

requ

ency

, f

= 40

(a)

(b)

(c )

T=798K

av = 35.3o

av = 27.5o

av = 32o

Fig. 6 Strain dependence of the distribution of misorientation angleof (sub)grain boundaries developed in 7075 Al alloy with pancakedgrains. The samples were deformed to various strains at 798 K and at2.9 × 10−3 s−1.

grains are subdivided and fine grains are evolved locally inseveral places at ε = 40% and in the almost whole area atε = 80%. It was difficult to carry out more detailed observa-tion of these SEM micrographs because the specimen surfacewas more corrugated and also contaminated by further defor-mation.

Figure 8 shows a series of OIM micrographs for the sameplace of the same specimen as in Fig. 7. OIM scanning wasperformed in a area of 68 µm × 46 µm with a step size of0.38 µm in a hexagonal grid. The colors and lines indicatethe same meaning as in Fig. 4. The initial layered pancakedgrains with many LABs (Fig. 8(a)) are subdivided by new

HABs and fine grains are inhomogeneously evolved at severalgrain interiors at ε = 40%. With increasing strain, new grainsare produced in colony and the number increases. Thesefine grains are rapidly developed after deformation to strainsabove 80%, leading to almost full development of a new grainstructure at strains higher than around 160% (Fig. 4).

The point-to-point misorientation (θ ), namely a relativedifference of crystal orientation between two adjacent scanpoints (0.38 µm), were measured along the line marked as A-B in a same pancaked grain interior in Fig. 8(a). Figure 9shows changes in misorientation angles measured in the ex-actly same position marked as A-B in Fig. 8 with variousstages of hot deformation. It can be seen in Fig. 9(a) thatsubgrain boundaries with a few degrees of the misorientationare developed in the pancaked grain interior at ε = 0. Withincreasing strain, a few HABs are evolved near the grain cor-ners and misorientations beyond 10◦ are also formed at sev-eral places in the grain interior (Fig. 9(b)). It can be seenin Fig. 9(c) that, with further straining to 80%, HABs withmisorientations ranging from 20◦ to beyond 40◦ are homoge-neously developed in the same position in the pancked grainin Fig. 8(a). This indicates that LABs developed in Fig. 9(a)are almost fully changed to HABs in Fig. 9(c) by straining to80%, although Fig. 9 is one of the typical results obtained.

4. Discussion

4.1 Superplasticity and grain boundary slidingIt is shown in the present experimental results described

above that the peak flow stress (σp) for an unrecrystallizedcoarse-grained 7075 Al alloy shows different strain rate de-pendence in the three regions of strain rate and a maximumelongation as well as the stress exponent of around 1.7 ap-pears in the region II (Fig. 2). All of these results is similarto the characteristics of typical superplasticity for fine grainedmaterials.1, 2) It is considered, therefore, that hot deformationof the present 7075 Al alloy can be controlled mainly by grainboundary sliding (GBS) and so approximated by the follow-ing usual equation applicable for the superplasticity.

ε = K1 D−pσ np exp(−Q/RT) (1)

where K1, p, n and Q are experimental constants and the oth-ers have usual physical meanings. If the σp can be controlledby the grain size in the ST direction (DST), as discussed in de-tail in elsewhere,12) the grain size exponent (p) obtained from

2404 X. Yang, H. Miura and T. Sakai

Fig. 7 SEM micrographs of the exactly same place in a L-ST plane of thesame sample deformed at 798 K and at 2.9 × 10−3 s−1. Tensile axis wasparallel to the layered grain boundaries. (a) ε = 10%, (b) ε = 40% and(c) ε = 80%.

the result in Fig. 3 is about 2.5. This value is almost similarto the reported values, i.e. p = 2 to 3.1, 2) The apparent ac-tivation energy for hot deformation (Q) is obtained as about138 kJ/mol from the results in Fig. 2. This value is roughlythe same as that for self diffusion of aluminum.15)

A SEM micrograph in the L-ST section of a specimenstrained to 40% at 798 K and at 2.9 × 10−3 s−1 is shown inFig. 10. It can be seen in Fig. 10 that scratched marker linesslip off across serrated layered HABs and also some of themare rotated especially in finer grains. Some portions of the ser-

Fig. 8 OIM micrographs in the same place as that in Fig. 7 (L-STplane). The samples were deformed to various strains at 798 K and at2.9 × 10−3 s−1. The line marked as A-B in (b) and (c) shows the exactlysame position as that in an originally pancaked grain in (a). The colors andlines have the meanings as described in Fig. 4. (a) ε = 0%, (b) ε = 40%and (c) ε = 80%.

rated HABs are locally inclined to the tensile axis, althoughthey are generally parallel to the latter (see also Figs. 7 and 8).It is easily understandable that GBS may locally take place atsuch serrated regions of rugged HABs.11, 12, 16)

The rotation angle (θrot) of such marker lines scratched per-pendicular to the tensile axis were measured by using someSEM micrographs as in Fig. 10. Changes in the rotation an-gle with deformation are shown in Fig. 11. The average valueof θrot and the scattering starts to increases rapidly just af-

Continuous Dynamic Recrystallization in a Superplastic 7075 Aluminum Alloy 2405

0

10

20

30

40

50

0

5

10

15

20

0

10

20

30

40

50

7075 Al M

isor

ient

atio

n an

gle,

/ d

eg

Distance, d / m

= 0

= 80

=40

A B

= 2.9x10-3s

-1

(a)

(c)

(b)

.T = 798K

10

Fig. 9 Misorientation distributions of (sub)grain and dislocation bound-aries developed in a pancaked grain interior measured along the linemarked as A-B indicated in Fig. 8. The samples were deformed to var-ious strains at 798 K and at 2.9 × 10−3 s−1.

ter the peak strain (εp∼= 5%) and approaches to a saturated

value, although the scattering becomes large. It is concludedthat GBS followed by grain rotation takes place frequentlyin layered HABs during hot deformation, and can acceleratedynamic development of new fine grains in pancaked grainstructures.

It is found from the results mentioned above that GBScan play an important role not only on the appearance ofsuperplasticity, but also on the dynamic evolution of newfine grains, that will be discussed in detail later. Such phe-nomena appearing during superplastic deformation have beentermed as continuous dynamic recrystallization (CDRX).1–7)

CDRX is different from conventional, i.e. discontinuous,DRX (DDRX) which is composed of the two stages, i.e. nu-cleation and growth of new grains.17) The details and themechanisms of CDRX have not been, however, fully unre-solved. This will be discussed in detail in the following sec-tion.

4.2 Continuous dynamic recrystallization (CDRX)From the present experimental results and the argument de-

scribed above, the process of new grain evolution occurringin an unrecrystallized pancaked grain structure can be shownschematically in Fig. 12. The layered grain boundaries devel-oped by prior cold-rolling are serrated by a progress of staticrecovery during reheating to high temperatures.10–12) First, at

Fig. 10 A SEM micrograph in the L-ST plane of 7075 Al alloy deformed toa strain of 40% at 798 K and at 2.9×10−3 s−1. Tensile axis was parallel tothe rolling direction. Arrowheads indicate a marker line scratched beforehot deformation. Note not only the operation of grain boundary sliding,but also grain rotation taking place in finer grains.

Nominal strain, 0 50 100 150 200

0

10

20

30

40

50

Rot

atio

n an

gle,

ro

t / d

eg

T = 798K = 2.9x10-3s-1

7075 Al

Fig. 11 Strain dependence of rotation angle of scratched marker lines per-pendicular to the tensile axis of 7075 Al alloy deformed at 798 K and at2.9 × 10−3 s−1.

an early stage of hot deformation, GBS takes place at sev-eral local places such as triple junctions and serrated layeredHABs, leading to the rotation of subgrains just beside theseHABs. The subgrain rotation should bring about increase inthe misorentation angles between subgrains beside the HABs.Then, GBS can operate even on these (sub)grain boundaries,

2406 X. Yang, H. Miura and T. Sakai

Fig. 12 Schematic illustration for new grain evolution taking place in alayered pancaked structure just after hot deformation is initiated (L-STplane). Thick, middle and thin lines indicate high, medium and low angle(sub)grain boundaries, respectively.

leading to the further rotation of neighboring subgrains andthen the evolution of new grain boundaries with medium tohigh anlge misorientations. Such a process can repeatedlyoperate in all area of the layered structure, finally followed bya full evolution of new grains with HABs. It is concluded,therefore, that this is a kind of deformation-induced continu-ous reaction, that is continuous DRX (CDRX).

The present model for dynamic evolution of new grainsis strongly supported by the several recent works.11, 12, 18–20)

They clarified the possible operation of GBS along the orig-inal HABs which can promote the subgrain rotation and therandomization of the initial strong texture accompanied byhot deformation. In the previous studies on CDRX tak-ing place during hot deformation of Al alloys, most ofthe tested materials had unrecrystallized and pan-caked lay-ered grain structures developed by prior cold- or warm-working.1–12),7–24) It may be concluded, generally that finegrain evolution due to CDRX in unrecrystallized Al alloyscan be explained by the present model shown in Fig. 12.

Previous models proposed for CDRX during superplasticdeformation explain how recovered subgrains developed ingrain interiors can transform in-situ to new grains with HABs.These models are briefly described as follows; Bricknell andEdington21) explained that new grains with HABs are evolvedby subgrain boundary dissociation and coalescence in graininteriors. Nes4, 22) proposed that HABs could be developedby the growth of subgrains in a sufficient distance if the mis-orientation angles can accumulate linearly with their migra-tion distance. Several other models for CDRX have been pro-posed and discussed, e.g. dislocations produced under defor-mation can be absorbed into subboundaries, resulting in anincrease in their misorientations,23, 24) and subgrain boundarysliding may cause their rotation,6) although it takes hardly inLABs. These models must assume that new grain evolutiontakes place homogeneously in all grain interiors. All of thesemodels for CDRX can explain how subgrains with LABs de-veloped can transform in-situ to new grains with HABs under

0 100 200 300 400 500 600 700

25

30

35

40

Ave

rage

mis

orie

ntat

ion,

av

/ de

g

Nominal strain,

7075 Al

T = 798K = 2.9x10-3s-1

Fig. 13 Strain dependence of average misorientation of (sub)grain bound-aries of 7075 Al alloy deformed at 798 K and at 2.9 × 10−3 s−1.

hot deformation, and so must assume that such phenomenatake place homogeneously in the whole matrices.

In contrast, the present model shown in Fig. 12 suggeststhat GBS starts to take place inhomogeneously at serratedHABs or severely declined ones against the tensile axis whichare developed by prior-cold or warm working followed by re-heating. GBS takes place locally along some portions of suchlayered HABs and subsequently in grain interiors under hotdeformation, as discussed above. It is concluded, therefore,that the evolution of new fine grains due to CDRX can re-sult from the operation of GBS initially at some local placessuch as the triple point junctions and the portions of serratedHABs, leading to the rotation of subgrains in these regions.With further deformation, GBS should be able to operate onthese new HABs evolved near original grain boundaries, fi-nally followed by full development of new grains in origi-nal grain interiors by subsequent GBS and grain rotation. Itis concluded, therefore, that pancaked grain structures intro-duced by cold rolling is an important prerequisition not onlyfor the appearance of superplasticity, but also for the dynamicevolution of new fine grains.

It is recently reported16, 25, 26) that new fine-grains are de-veloped by grain subdivision due to development of defor-mation bands under large strain deformation even at low tointermediate temperature (T < 0.5Tm, where Tm is the melt-ing point). The number and the average misorientions (θav) ofthese dislocation boundaries rapidly rise with deformation, fi-nally leading to development of a new grained structure. Thisphenomena is also called as CDRX.25, 26) In this case the θav

increase from zero to a saturation value of 25–30◦ in severelarge strain (e.g. ε = 104–105%).25) In contrast, CDRX op-erating at around strains below 102 to 103% under hot defor-mation is not the same as that under low temperature largestrain deformation. Figure 13 shows strain dependence ofthe θav for (sub)grain boundaries obtained from Fig. 6 for thepresent 7075 Al alloy. Here the θav starts to increase rapidlyfrom about 27.5◦ just before deformation, rises at a steadyrate against strain at strains of above 160%, and approachesto around 40◦ in high strain. This θav − ε behavior for theunrecrystallized 7075 Al alloy can be caused by the follow-ing reasons; high density layered HABs developed by prior

Continuous Dynamic Recrystallization in a Superplastic 7075 Aluminum Alloy 2407

cold-rolling exist stably before deformation and GBS takesplace frequently on these HABs, leading to rapid increase inθav during early deformation at T > 0.5Tm.

5. Conclusions

New grain evolution as well as superplasticity was stud-ied in a prior cold-rolled 7075 Al alloy with an unrecrystal-lized pancaked grain structure and the main results obtainedare summarized as follows.

(1) Strain rate dependence of flow stress (and total elon-gation to fracture) is sensitively changed in the three regionsof strain rate and affected by prior cold-rolling reduction.Typical superplasticity takes place in the region of mediumstrain rate.

(2) Prior cold-rolling develops a layered, pancaked grainstructure and so brings about high density layered high-anglegrain boundaries per uint volume. Such pancaked grain struc-tures are an important prerequisition not only for dynamicevolution of new grains, but also for the appearance of su-perplasticity.

(3) Grain boundary sliding (GBS) takes place at someportions of original layered grain boundaries, leading to rota-tion of subgrains just beside the latters at early stages of defor-mation. Subsequently GBS can operate even at the (sub)grainboundaries with medium to high angle misorientions evolvednear grain boundaries. These processes takes place repeat-edly in the whole matrices, finally leading to full evolution ofa fine grain structure.

(4) It is concluded that dynamic evolution of fine grainsas well as superplasticity in the unrecrystallized 7075 alu-minum alloy can be controlled manly by GBS and so thestructural mechanism is a deformation-induced continuous re-action, that is continuous dynamic recrystallization.

Acknowledgments

The authors are intended to the following bodies for the fi-nancial support: Ministry of Education, Science and Culture(Grant-in-Aid for Scientific Research (c) (No. 12650710))and the Light Metals Educational Foundation, Japan.

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III, 26, (Springel-Verlag, Berlin, 1990) pp. 647.16) It should be noted that GBS takes place on HABs in not only the L-ST

plane, but also the L-LT plane. The initial grain size of about 17 µmis reduced to about 5.7 µm in the ST direction and elongated to about51 µm along the L (i.e. rolling) direction by a reduction of 67%, leadingto formation of a pan-caked layered grain structure (Figs. 4 and 5). Thedensity of HABs is significantly low in the L-LT plane compared withthat in the L-ST plane. Then the role of GBS in the L-LT plane can benegligible small compared with that in the L-ST plane. It was discussedin elsewhere,11, 12) based on the effect of prior cold rolling reductionon hot deformation, that superplasticity is dependent sensitively on thegrain size in the ST direction.

17) F. J. Humphreys and M. Hatherly: Recrystallization and Related An-nealing Phenomena, (Pergamon, 1995) pp. 19–38, pp. 167–171.

18) H. Gudmundsson, D. Brooks and J. A. Wert: Acta Metall. Mater. 39(1991) 19–35.

19) T. R. McNelley and M. E. McMahon: Metall. Trans. A 27A (1996)2252–2262.

20) J. Liu and D. J. Chakrabart: Acta Metall. 44 (1996) 4647–4661.21) R. H. Bricknel and J. W. Edington: Metall. Trans. A 10A (1979) 1257–

1263.22) E. Nes: Superplasticity, ed. by B. Baudelet and M. Surey, Center

National de 1a Rechaerche Scintifique, (Paris, France, 1985) pp. 7.1–7.14.

23) Q. Liu, X. Huang, M. Yao and J. Yang: Acta Metall. 40 (1992) 1753–1762.

24) K. Tsuzaki, X. Huang and T. Maki: Acta Metall. 44 (1996) 4491–4499.25) T. Sakai, A. Belyakov and H. Miura: Recrystallization and Grain

Growth, eds. by G. Gottstein and D. A. Molodov (Springer, 2001)pp. 669–682.

26) A. Belyakov, T. Sakai, H. Miura and K. Tsuzaki: Philos. Mag. A 81(2001) 2629–2643.